481 Silicon–Germa 22. Silicon–Germanium: Properties, Growth and Applications Silicon–germanium is an important material that is used for the fabrication of SiGe heterojunction bipolar transistors and strained Si metal–oxide– semiconductor (MOS) transistors for advanced complementary metal—oxide–semiconductor (CMOS) and BiCMOS (bipolar CMOS) technologies. It also has interesting optical properties that are increasingly being applied in silicon-based photonic devices. The key benefit of silicon–germanium is its use in combination with silicon to produce a heterojunction. Strain is incorporated into the silicon–germanium or the silicon during growth, which also gives improved physical properties such as higher values of mobility. This chapter reviews the properties of silicon–germanium, beginning with the electronic properties and then progressing to the optical properties. The growth of silicon–germanium is considered, with particular emphasis on the chemical vapour deposition technique and selective epitaxy. Finally, the properties of polycrystalline silicon–germanium are discussed in the context of its use as a gate material for MOS transistors. 22.1 22.1.3 22.1.4 22.1.5 Dielectric Constant .................... Density of States ....................... Majority-Carrier Mobility in Strained Si1−x Gex .................. 22.1.6 Majority-Carrier Mobility in Tensile-Strained Si on Relaxed Si1−x Gex .................. 22.1.7 Minority-Carrier Mobility in Strained Si1−x Gex .................. 22.1.8 Apparent Band-Gap Narrowing in Si1−x Gex HBTs........................ 484 484 486 486 486 487 22.2 Optical Properties of SiGe ..................... 22.2.1 Dielectric Functions and Interband Transitions .......... 22.2.2 Photoluminescence ................... 22.2.3 SiGe Quantum Wells .................. 488 22.3 Growth of Silicon–Germanium .............. 22.3.1 In-Situ Hydrogen Bake .............. 22.3.2 Hydrogen Passivation ................ 22.3.3 Ultra-Clean Epitaxy Systems ....... 22.3.4 Si1−x Gex Epitaxy ........................ 22.3.5 Selective Si1−x Gex Epitaxy........... 492 492 492 492 492 492 488 489 490 References .................................................. 497 Silicon–germanium (Si1−x Gex ) alloys have been researched since the late 1950s [22.1], but it is only in the past 15 years or so that these layers have been applied to new types of transistor technology. Si1−x Gex was first applied in bipolar technologies [22.2, 3], but more recently has been applied to metal–oxide–semiconductor (MOS) technologies [22.4–7]. This has been made possible by the development of new growth techniques, such as molecular-beam epitaxy (MBE), low-pressure chemical vapour deposition (LPCVD) and ultra-highvacuum chemical vapour deposition (UHV-CVD). The key feature of these techniques that has led to the development of Si1−x Gex transistors is the growth of epitaxial layers at low temperatures (500–700 ◦ C). This allows Si1−x Gex layers to be grown without disturbing the doping profiles of structures already present in the silicon wafer. Si1−x Gex layers can be successfully grown on silicon substrates even though there is a lattice mismatch between silicon and germanium of 4.2%. The primary property of Si1−x Gex that is of interest for bipolar transistors is the band gap, which is smaller than that of silicon and controllable by varying the germanium content. Band-gap engineering concepts, which were previously only possible in compound semiconductor technologies, have now become viable in silicon technology. These concepts have introduced new degrees of freedom in the design of bipolar transistors that have led to dramatic improvements in transistor Part C 22 22.4 Polycrystalline Silicon–Germanium........ 494 22.4.1 Electrical Properties of Polycrystalline Si1−x Gex .......... 496 Physical Properties of Silicon–Germanium ......................... 482 22.1.1 Critical Thickness ....................... 482 22.1.2 Band Structure.......................... 483 482 Part C Materials for Electronics performance. In Si1−x Gex heterojunction bipolar transistors (HBTs), the Si1−x Gex layer is incorporated into the base and the lower band gap of Si1−x Gex than Si is used to advantage to dramatically improve the high-frequency performance. Si1−x Gex HBTs have been produced with values of cut-off frequency, f T , approaching 300 GHz [22.8], a value unimaginable in silicon bipolar transistors. Values of gate delay well below 10 ps can be achieved in properly optimised Si1−x Gex HBTs [22.9]. In Si1−x Gex MOS field-effect transistors (MOSFETs), the Si1−x Gex layer is incorporated in the channel and is used to give improved values of mobility. Initially, strained Si1−x Gex layers on silicon substrates were used to give improved hole mobility in p-channel transistors [22.7], but more recently thin, strained silicon layers on relaxed SiGe virtual substrates have been used to give improvements in both electron and hole mobility [22.4–6]. In this chapter, the properties of single-crystal silicon–germanium will first be outlined, followed by a description of the methods used for growing silicon– germanium layers. The properties and applications of polycrystalline silicon–germanium are also discussed later in the article. 22.1 Physical Properties of Silicon–Germanium Silicon and germanium are completely miscible over the full range of compositions and hence can be combined to form Si1−x Gex alloys with the germanium content, x, ranging from 0 to 1 (0–100%). Si1−x Gex has a diamondlike lattice structure and the lattice constant is given by Vegard’s rule: aSi Ge = aSi + x aGe − aSi , (22.1) 1−x x Part C 22.1 where x is the germanium fraction and a is the lattice constant. The lattice constant of silicon, aSi , is 0.543 nm, the lattice constant of germanium, aGe , is 0.566 nm and the lattice mismatch is 4.2%. When a Si1−x Gex layer is grown on a silicon substrate, the lattice mismatch at the interface between the Si1−x Gex and the silicon has to be accommodated. This can either be done by compression of the Si1−x Gex layer so that it fits to the silicon lattice or by the creation of misfit dislocations at the interface. These two possibilities are illustrated schematically in Fig. 22.1. In the former case, the Si1−x Gex layer adopts the silicon lattice spacing in the plane of the growth and hence the normally cubic Si1−x Gex crystal is distorted. When Si1−x Gex growth occurs in this way, the Si1−x Gex layer is under compressive strain and the layer is described as pseudomorphic. In the second case, the Si1−x Gex layer is unstrained, or relaxed, and the lattice mismatch at the interface is accommodated by the formation of misfit dislocations. These misfit dislocations generally lie in the plane of the interface, as shown in Fig. 22.1, but dislocations can also thread vertically through the Si1−x Gex layer. Unstrained Si1–x Gex Silicon substrate Pseudomorphic Misfit dislocation Fig. 22.1 Schematic illustration of pseudomorphic Si1−x Gex growth and misfit-dislocation formation Silicon–Germanium: Properties, Growth and Applications 22.1.1 Critical Thickness particular importance to both Si1−x Gex HBTs and Si1−x Gex MOSFETs, is the effect of a silicon cap layer, which has been shown to increase the critical thickness of the underlying Si1−x Gex layer. Figure 22.3 shows a comparison of the calculated critical thickness as a function of germanium percentage for stable Si1−x Gex layers with and without a silicon cap. It can be seen that the critical thickness is more than doubled by the presence of the silicon cap. The presence of misfit dislocations in devices is highly undesirable, since they create generation/ recombination centres, which degrade leakage currents when they are present in the depletion regions of devices. Threading dislocations also highly undesirable, as they can lead to the formation of emitter/collector pipes in Si1−x Gex HBTs. When designing Si1−x Gex devices, it is important that the Si1−x Gex thickness is chosen to give a stable layer, so that dislocation formation is avoided. A base thickness below the silicon cap curve in Fig. 22.3 will ensure a stable layer, which will withstand ion implantation and high-temperature annealing without encountering problems of relaxation and misfit-dislocation generation. Considerable research has been done on the oxidation of Si1−x Gex [22.14, 15], and it has been found that the germanium in the Si1−x Gex layer does not oxidise, but piles up at the oxide/Si1−x Gex interface. This pile-up of germanium makes it difficult to achieve low values of interface state density in oxidised Si1−x Gex layers. It is therefore advisable to avoid direct oxidation of Si1−x Gex layers, particularly in Si1−x Gex MOS technologies. In Si1−x Gex MOSFETs, a silicon cap is often Critical thickness (nm) Critical thickness (nm) 1000 1000 Relaxed PeopleBean 100 Si cap 100 Metastable MathewsBlakeslee no Si cap 10 10 Stable 1 1 5 10 15 30 20 25 Germanium percentage (%) 5 10 15 30 20 25 Germanium percentage (%) Fig. 22.3 Critical thickness as a function of germanium Fig. 22.2 Critical Si1−x Gex thickness as a function of ger- manium percentage (after Iyer et al. [22.2], copyright 1989 IEEE) 483 Part C 22.1 There is a maximum thickness of Si1−x Gex that can be grown before relaxation of the strain occurs through the formation of misfit dislocations. This is known as the critical thickness of the Si1−x Gex layer, and depends strongly on the germanium content, as shown in Fig. 22.2. The original calculations of critical layer thickness were made by Matthews and Blakeslee [22.11, 12] on the basis of the mechanical equilibrium of an existing threading dislocation. However, measurements of dislocation densities in Si1−x Gex showed, in many cases, no evidence of misfit dislocations for Si1−x Gex layers considerably thicker than the Matthews–Blakeslee limit. These results were explained by People and Bean [22.13] who calculated the critical thickness on the assumption that misfitdislocation generation was determined solely by energy balance. The discrepancy between these two types of calculation can be explained by the observation that strain relaxation in Si1−x Gex layers occurs gradually. Layers above the People–Bean curve can be considered to be completely relaxed, whereas layers below the Matthews–Blakeslee curve can be considered to be fully strained. These fully strained layers are termed stable and will not relax during any subsequent high temperature processing. Layers lying between the two curves are termed metastable; these layers may be free of dislocations after growth, but are susceptible to relaxation during later high-temperature processing. In practice, a number of additional factors influence the critical thickness of a Si1−x Gex layer. Of 22.1 Physical Properties of Silicon–Germanium percentage for stable Si1−x Gex layers with and without a silicon cap (after Jain et al. [22.10], copyright 1992 Elsevier) 484 Part C Materials for Electronics included above the Si1−x Gex layer that can be oxidised to create the gate oxide. Energy (eV) 0.35 22.1.2 Band Structure Si1−x Gex alloys have a smaller band gap than silicon partly because of the larger lattice constant and partly because of the strain. Figure 22.4 shows the variation of band gap with germanium percentage for strained and unstrained Si1−x Gex . It can be seen that the strain has a dramatic effect on the band gap of Si1−x Gex . For 10% germanium, the reduction in band gap compared with silicon is 92 meV for strained Si1−x Gex , compared with 50 meV for unstrained Si1−x Gex . The variation of band gap with germanium content for strained Si1−x Gex can be described by the following empirical equation: E G (x) = 1.17 + 0.96x − 0.43x 2 + 0.17x 3 (22.2) The band alignment for compressively strained Si1−x Gex on unstrained silicon is illustrated schematically in Fig. 22.5. This band alignment is referred to as type I, and the majority of the band offset at the heteroBandgap (eV) 1.2 1.1 Unstrained 1.0 Part C 22.1 Strained 0.9 0.8 0 10 20 40 30 Germanium percentage (%) Fig. 22.4 Band gap as a function of germanium percentage for strained [22.16] and unstrained [22.1] Si1−x Gex (after Iyer et al. [22.2], copyright 1989 IEEE) Ec Ev Unstrained silicon ∆EG 0.30 Compressively strained Si1 – x Gex Fig. 22.5 Schematic illustration of the band alignment ob- tained for a compressively strained Si1−x Gex layer grown on an unstrained silicon substrate 0.25 ∆EV 0.20 0.15 0.10 ∆EC 0.05 0.00 0 10 40 20 30 Germanium percentage (%) Fig. 22.6 Valence- and conduction-band offsets as a function of germanium percentage for strained Si1−x Gex grown on unstrained silicon (after Poortmans et al. [22.17], copyright 1993 Elsevier) junction interface occurs in the valence band, with only a small offset in the conduction band. Different band alignments can be obtained by engineering the strain in the substrate and the grown layer in different ways. For example, type II band alignments can be obtained by growing tensile-strained silicon on top of unstrained Si1−x Gex . This arrangement gives large conductionand valence-band offsets and is used in strained silicon MOSFETs. Figure 22.6 shows the variation of valence-band offset, ∆E V , conduction-band offset, ∆E C , and band-gap narrowing, ∆E G , with germanium content. It can be seen that the majority of the band offset occurs in the valence band. For example for 10% germanium, the valenceband offset is 0.073 eV, compared with 0.019 eV for the conduction-band offset. The conduction-band offset can therefore be neglected for most practical purposes. 22.1.3 Dielectric Constant The dielectric constant of Si1−x Gex can be obtained by linear interpolation between the known values for silicon and germanium [22.17] using the following equation: ε(x) = 11.9(1 + 0.35x) (22.3) 22.1.4 Density of States While, the density of states in the conduction band in Si1−x Gex is generally assumed to be the same as that in silicon, there is some evidence in the literature to suggest that the density of states in the valence band Silicon–Germanium: Properties, Growth and Applications Ev – EF (eV) 20 % 0.06 30 % 0% 10 % 0.05 22.1 Physical Properties of Silicon–Germanium 485 These results are shown in Fig. 22.7 for Si1−x Gex with four different germanium contents. It can be seen that the Fermi level moves deeper into the valence band as the germanium concentration increases. Figure 22.8 shows the ratio of the calculated density of states in the valence 0.04 a) Electron mobility (103 cm2 / Vs) 0.03 3 Si1–xGex on <001> Si 300K 0.02 0.01 0.00 1018 1019 1020 Hole concentration (cm–3) 2 µ⊥ Fig. 22.7 Fermi-level position as a function of hole con- centration for Si1−x Gex with four different germanium concentrations (after Iyer et al. [22.2], copyright 1989 IEEE) No alloy scattering µ⏐⏐ 1 µ⊥ NV(Si1–xGex)/NV(Si) 1.2 µ⏐⏐ 1.0 0 0.0 0.8 0.6 0.2 0.4 0.6 Alloy scattering 1.0 0.8 Ge mole fraction b) Hole mobility (102 cm2 / Vs) 0.4 3 0.2 2 0.0 0 5 10 15 Fig. 22.8 Ratio of density of states in the valence band for Si1−x Gex to that for Si as a function of germanium percentage (after Poortmans [22.20], copyright 1993 University of Leuven) 1049 8 7 6 5 No alloy scattering 4 µ⊥ 3 2 is considerably smaller. Manku and Nathan [22.18, 19] have calculated the E–k diagram for strained Si1−x Gex and shown that the density-of-states hole mass is significantly lower, by a factor of approximately three at 30% germanium. There is some experimental evidence to support this calculation. For example, freeze-out of holes in p-type Si1−x Gex has been reported to occur at higher temperatures than in p-type silicon [22.20] and enhancements in the majority-carrier, hole mobility have been reported for p-type Si1−x Gex [22.21]. Using the calculated values of hole density of states of Manku and Nathan [22.18,19], the hole concentration can be calculated as a function of Fermi-level position. µ⏐⏐ Part C 22.1 35 20 25 30 Germanium percentage (%) µ⏐⏐ 1039 8 7 6 5 µ⊥ Alloy scattering 4 3 0.0 0.2 0.4 0.6 1.0 0.8 Ge mole fraction Fig. 22.9a,b Calculated 300-K electron (a) and hole (b) in- plane and out-of-plane low-field mobilities in strained Si1−x Gex grown on (100) Si (after Fischetti et al. [22.22], copyright 1996 American Institute of Physics) 486 Part C Materials for Electronics band for Si1−x Gex to that for silicon as a function of germanium content. It is clear that the density of states in the valence band for Si1−x Gex is significantly lower than that for silicon at germanium contents of practical interest. 22.1.5 Majority-Carrier Mobility in Strained Si1−x Gex Values of in-plane and out-of-plane low-field mobility in strained Si1−x Gex grown on (100) Si have been calculated by Fischetti et al. [22.22], and are shown in Fig. 22.9. There is some uncertainty in the chosen values of alloy scattering parameters used in the calculations, but nevertheless the results are representative of current understanding. These results show a large enhancement of low-field hole mobility for Si1−x Gex compared with unstrained silicon, but only a modest enhancement of low-field electron mobility. These results indicate that strained Si1−x Gex channels can be used to significantly improve the mobility of p-channel MOSFETs, but little benefit is obtained for n-channel MOSFETs. For this reason, industry focus has moved away from channels realised in Si1−x Gex to channels realised in tensile-strained silicon, as discussed below. 22.1.6 Majority-Carrier Mobility in Tensile-Strained Si on Relaxed Si1−x Gex Part C 22.1 Tensile-strained silicon can be produced by growing a thin silicon layer on top of a relaxed Si1−x Gex virtual substrate. Figure 22.10 shows a typical virtual substrate for a surface-channel MOS transistor. A graded Si1−x Gex layer is grown on top of the silicon substrate with the Ge content varying from 0 to 30%. Misfit dislocations will form in this layer, but the majority of dislocations will be in the plane of the Si1−x Gex layer and only a small percentage will propagate vertically Effective mobility, µeff (cm2/Vs) 1600 Buried strained-Si 1400 1200 Surface strained-Si 1000 800 600 400 200 0 0 CZ Si control VDrain = 40 mV 290 K 0.1 0.2 0.3 0.4 0.5 0.6 Effective electric field, Eeff (MV/cm) Fig. 22.11 Effective electron mobility as a function of effective electric field for strained Si MOSFETs fabricated on a 30% Si1−x Gex virtual substrate (after Welser et al. [22.23], copyright 1994 IEEE) to the surface of the layer. A 30%-relaxed Si1−x Gex buffer is then grown followed by a thin tensile-strained Si1−x Gex layer in which the channel is fabricated. The key to any virtual substrate growth process is the minimisation of dislocation propagation to the surface of the wafer. Figure 22.11 shows typical values of effective electron mobility obtained from measurements on n-channel MOS transistors for a Si1−x Gex virtual substrate with 30% Ge [22.23]. For the surface-channel strained Si device, the effective mobility is enhanced by 80% compared with the Si control transistor due to the tensile strain in the Si1−x Gex layer. Enhanced hole mobility can also be obtained in tensile-strained Si grown on a Si1−x Gex virtual substrate, though higher germanium contents are needed to obtain a significant mobility enhancement. Figure 22.12 shows typical values of effective hole mobility in strained Si for Ge contents between 35 and 50% [22.24]. The effective mobility of the strained Si device is enhanced by 100% compared with the Si control device. Strained Si Relaxed Si0.7Ge0.3 Graded Si1 – xGex, 0 – 30 % Fig. 22.10 Schematic illustration of a typical tensilestrained Si layer grown on top of a Si1−x Gex virtual substrate 22.1.7 Minority-Carrier Mobility in Strained Si1−x Gex Si1−x Gex HBTs are minority-carrier devices and hence values of the minority-carrier mobility of more interest than the majority-carrier mobility. Unfortunately very few measurements of minority-carrier mobility have been made in Si1−x Gex . Poortmans [22.20] inferred values of minority-carrier mobility from measure- Silicon–Germanium: Properties, Growth and Applications Effective mobility (cm2/Vs) Bulk Si 35 % Ge 40% Ge 45 % Ge 50% Ge 250 200 150 100 50 0 0.2 0.3 0.4 0.5 0.8 0.6 0.7 Effective field (MV/cm) Fig. 22.12 Effective hole mobility as a function of effective electric field for strained Si MOSFETs fabricated on a Si1−x Gex virtual substrate with Ge contents in the range 35–50% (after Leitz et al. [22.24], copyright 2002 IEEE) NCNVDnb ratio 0.45 0.40 0.35 8% 0.30 22.1 Physical Properties of Silicon–Germanium 487 data directly obtained from measurements on Si1−x Gex HBTs. The gain enhancement in a Si1−x Gex HBT is determined by the ratio of the product NC NV Dnb in Si1−x Gex and Si, together with the band-gap narrowing due to the strained Si1−x Gex . Figure 22.13 shows a graph of this NC NV Dnb ratio as a function of acceptor concentration for three values of germanium content. It can be seen that for germanium contents of practical interest, in the range 11–16%, this ratio has a value of around 0.25. 22.1.8 Apparent Band-Gap Narrowing in Si1−x Gex HBTs In Si1−x Gex HBTs, the apparent band-gap narrowing is often quoted, which combines the effect of the bandgap reduction and the effect of high doping. Poortmans et al. [22.17] have developed a theoretical approach that has been shown to be in reasonable agreement with experiment. Figure 22.14 shows the apparent band-gap narrowing in Si1−x Gex as a function of acceptor concentration for three values of germanium content. At low acceptor concentrations, the apparent band-gap narrowing in Si1−x Gex is slightly higher than that in silicon, but at acceptor concentrations in the range 1–2 × 1019 cm−3 , the apparent band-gap narrowing is approximately the same. This latter doping range is the base doping range that is of practical interest for Si1−x Gex HBTs. 12 % 0.25 16 % 10 –19 10 –20 Acceptor concentration (cm–3) Fig. 22.13 Measured values of the ratio of NC NV Dnb in Si1−x Gex to that in Si as a function of acceptor concentration (after Poortmans [22.20], copyright 1993 University of Leuven) ments on Si1−x Gex HBTs and found an enhancement in mobility compared with silicon by a factor of 1.2–1.4 for base doping concentrations in the range 5 × 1018 –5 × 1019 cm−3 . Given the scarcity of measured data on minority-carrier mobility and density of states in Si1−x Gex , the most reliable way of calculating the expected gain improvement in a Si1−x Gex HBT is to use 0.08 0.07 10 % 20 % 0.06 0.05 0% 0.04 0.03 18 10 1019 Acceptor concentration (cm–3) Fig. 22.14 Apparent band-gap narrowing as a function of acceptor concentration for Si1−x Gex with three different germanium concentrations (after Poortmans et al. [22.17], copyright 1993 Elsevier) Part C 22.1 0.20 0 Apparent bandgap narrowing (eV) 488 Part C Materials for Electronics 22.2 Optical Properties of SiGe Interest in the optical properties of SiGe stems from the desire to design silicon-based optoelectronic devices as well as the usefulness of many optical techniques in material analysis. The optical properties of bulk SiGe provides an important starting point for any attempt at in-depth understanding, however, nearly all real applications of SiGe involve the use of thin, strained layers. Beyond this, the formation of SiGe quantum structures within silicon devices has remained one of the most promising methods by which device engineers hope to improve the largely unimpressive optical behaviour of silicon that is brought about by its indirect band gap. The growth of quantum wells, quantum wires and quantum dots in the Si/SiGe system has been extensively explored [22.26–31]. However, there has as yet been little success at using Si/SiGe quantum structures to produce efficient silicon-based light emitters, although there have been impressive attempts [22.31]. Perhaps the most promising devices based on SiGe quantum wells are near-infra-red photodetectors in which the SiGe can be used to enhance sensitivity at the optical-communications wavelengths [22.32–36]. More futuristic applications of SiGe quantum wells include devices based on transitions between the confined energy levels in quantum wells. Devices based on these inter-subband transitions include quantum-well infrared photodetectors (QWIPs) [22.37, 38] and quantum cascade lasers [22.39, 40]. 22.2.1 Dielectric Functions and Interband Transitions A range of 1-µm-thick Si1−x Gex films grown by MBE on Si(100) substrates have been studied by spectroscopic ellipsometry to yield their complex dielectric functions at room temperature [22.25]. Both the real (ε1 ) and imaginary (ε2 ) parts of the measured dielectric function for Ge compositions of 0, 0.2, 0.4, 0.6, 0.8 and 1.0 are shown in Fig. 22.15a and b, respectively. In Fig. 22.15b the absorption structures observed at 3.4 and 4.2 eV in the spectra shown for silicon originate from direct band-to-band transitions at various regions in the Brillouin zone of silicon. The structure seen around 3.4 eV is due to E 0 , E 1 and E 1 + ∆1 interband transition b) ε2 a) ε1 60 Si1–xGex x x 40 x 50 = 0.0 Si1–xGex = 0.2 40 = 0.4 Part C 22.2 30 20 x = 0.6 x = 1.0 x 20 = 0.8 x 0 x x = 0.2 = 0.4 = 0.6 x = 0.8 x x 10 –20 2 3 4 5 Photon energy (eV) 0 2 = 0.0 = 1.0 3 4 5 Photon energy (eV) Fig. 22.15a,b Real (a) and imaginary (b) parts of the dielectric functions of relaxed Si1−x Gex alloys with composition x indicated in the legend (after Bahng et al. [22.25], copyright 2001, American Physical Society) Silicon–Germanium: Properties, Growth and Applications Photon energy (eV) 4.6 4.4 4.2 3.4 3.2 3.0 2.8 E0' E1 E0 + ∆1 E2(X) E2(Σ) Quadratic fit 2.6 2.4 2.2 2.0 1.8 0.0 0.2 0.4 0.6 1.0 0.8 Ge composition (x) 22.2 Optical Properties of SiGe 489 photoluminescence spectra is nearly always sufficient to allow a simple qualitative assessment of material quality; alternatively, detailed analysis can permit a broad range of material or structural parameters to be assessed or determined. No two photoluminescence spectra are the same. In this section and the section that follows on photoluminescence studies of Si/SiSe quantum wells, we will present a range of spectra that represent most of the key features that have been observed. Weber and Alonso [22.41] have provided a very useful study of the near-band-gap photoluminescence of bulk SiGe alloys. In their study bulk SiGe samples are cut from nominally undoped polycrystalline ingots prepared by a zone-levelling technique. Figure 22.17 shows the photoluminescence spectra for a range of compositions. Samples were excited using the 514-nm line of an argon ion laser and the sample temperature was 4.2 K. Fig. 22.16 Evolution of E 0 , E 1 , E 1 + ∆1 , E 2 (X) and E 2 (Σ) transition energies for relaxed Si1−x Gex with composition x (after Bahng et al. [22.25], copyright 2001, American Physical Society) E 1 (x) = 3.398 − 1.586x + 0.27x (22.5) E 1 + ∆1 (x) = 3.432 − 1.185x + 0.065x 2 (22.6) E 2 (X)(x) = 4.259 − 0.052x + 0.084x 2 (22.7) E 2 (Σ)(x) = 4.473 − 0.139x + 0.072x (22.8) 2 Si0.20Ge0.80 TO XGe – Ge TO XSi– Ge TO XSi – Si Si0.42Ge0.58 TO XGe – Ge XNP TO XSi– Ge XTA TO XSi – Si TO XGe – Ge Si0.62Ge0.38 XNP 22.2.2 Photoluminescence The emission properties of semiconductor structures are of fundamental interest to scientists as well as being an important analytical technique for engineers. In general, the features of low-temperature photoluminescence spectra are very dependent on the specific conditions under which materials are grown and treated. This is because emission energies and emission rates are often sensitive to even small variations of impurity or defect densities, as well as variations in strain or composition. A brief examination of low-temperature XTA XNP Part C 22.2 (22.4) 2 XTA XLA TO XGe – Ge TO XSi –Ge edges, whereas the structure at 4.2 eV is due to E 2 (X) and E 2 (Σ) edges [22.25]. The evolution of each of these transition edges for the full range of SiGe compositions is shown in Fig. 22.16. The quadratic fits shown in Fig. 22.16 are as follows [22.25]: E 0 (x) = 3.337 − 0.348x + 0.222x 2 XNP Si0.07Ge0.93 TO XSi– Ge XTA TO XSi – Si TO XSi – Si XNP Si0.92Ge0.08 XTA 1.15 1.05 0.95 0.85 Energy (eV) Fig. 22.17 Near-band-gap photoluminescence spectra for several bulk SiGe samples (after Weber et al. [22.41], copyright 1989, American Physical Society) 490 Part C Materials for Electronics Excitonic emission lines are a strong feature of photoluminescence spectra when the thermal energy of the semiconductor is less than the exciton binding energy. Each spectrum featured in Fig. 22.17, across the full range of SiGe compositions, show similar excitonic features. In most spectra the most pronounced peak is the nophonon (XNP ) line caused by the optical recombination of excitons bound to shallow impurities. In the case of the no-phonon line, momentum is conserved through interaction with the binding impurity. There are many candidate atoms for these shallow impurities and with B, P and As having binding energies of 4.2, 5.0 and 5.6 meV, respectively [22.42]. These bound energy states will tend to dominate luminescence spectra for doped samples and will always tend to be present in nominally undoped samples. The no-phonon line is accompanied by transverse-optical (XTO ) or transverse acoustic (XTA ) phonon replicas that are created as photon emission is accompanied by the momentum-conserving creation of lattice vibrations in Si–Si, Si–Ge or Ge–Ge bonds. Figure 22.18 shows how the spectrum from a bulk Si0.915 Ge0.085 sample develops with increasing temperature [22.41]. With increasing temperature the XNP line thermalises to leave the free-exciton (FENP ) line. Here, emission from nominally free excitons is greatly enhanced for the alloy samples by local fluctuations in composition that provide momentum-conserving scattering centres [22.43]. Part C 22.2 Luminescence intensity 1.14 1.12 1.10 1.08 Si1–xGex : x = 0.085 1.06 Energy (eV) 1.04 XTO Si1–xGex : x = 0.58 25 K FENP XNP x 10 LNP x 1.3 14 K x2 9K x 1.8 XTA 6K 1.22 1.26 XTO Ge – Ge XTO Ge – Si 1.26 XTO Si – Si 1.26 x1 1.26 Wavelength Fig. 22.19 Photoluminescence spectra of a bulk Si0.42 Ge0.58 sample at different temperatures (after Weber et al. [22.41], copyright 1989, American Physical Society) At higher temperatures the free-exciton line is also thermalised and all fine structure is lost, at temperatures around 25 K broad luminescence bands are commonly observed (Fig. 22.19 [22.41]. These deep luminescence bands are difficult to assign and have been ascribed to impurities, structural defects and, as in the case of the line presented in Fig. 22.19, potential wells formed by alloy fluctuations [22.41]. Weber and Alonso [22.41] use their data to provide analytical expressions for both the XNP and L bands for bulk Si1−x Gex in the range 0 ≤ x ≤ 0.85 as follows: L E gx (x) = 2.010 − 1.270x eV XNP 15 K 0.92 (x) E gx (x) = 1.155 − 0.43x + 0.206x 2 eV FETO NP FE Energy (eV) Luminescence intensity 1.00 0.96 (22.9) (22.10) x2 9K x2 6K x1 At low temperatures narrow excitonic luminescence features are indicative of defect-free material, and in this way low-temperature photoluminescence becomes a good qualitative tool with which material quality can be assessed. 4.2 K x1 22.2.3 SiGe Quantum Wells 1.08 1.10 1.12 1.14 1.16 1.20 1.18 Wavelength (µm) Fig. 22.18 Photoluminescence spectra of a bulk Si0.915 Ge0.085 sample at different temperatures (after Weber et al. [22.41], copyright 1989, American Physical Society) Figure 22.20 shows the first excitonic luminescence spectra from a Si/SiGe multiple quantum well grown by atmospheric-pressure CVD [22.28]. As we can see, many of the features seen in the photoluminescence spectra of quantum-well samples are similar to Silicon–Germanium: Properties, Growth and Applications Fig. 22.20 Excitonic photoluminescence spectrum of a SiGe quantum well (after Grutzmacher et al. [22.28], copyright 1993, AVS) those seen from the bulk samples described in the previous section. Again, the most pronounced peak is the no-phonon (NP) line and this is accompanied by phonon replicas, including, impressively, a twophonon replica of the NP line TO + TOSi−Si . The most significant difference between bulk and quantum well spectra is the energy positions of the excitonic features as these are shifted by quantum confinement effects. Robbins et al. [22.26], have provided one of the most detailed studies of near-band-gap photoluminescence from pseudomorphic SiGe layers and provide analytical expressions for all factors pertaining to the energy positions of the excitonic energy gap for Si1−x Gex quantum wells in the range 0 < x < 0.24. The effects of alloying, confinement, band offsets, alignment type and exciton binding energy are all taken into account. Samples used in the study were grown by low-pressure CVD at 920 ◦ C; a typical set of photoluminescence spectra are shown in Fig. 22.21. The exciton band gap at 4.2 K is considered for thick S ) where the energies are not (50-nm) strained layers (E X affected by quantum shifts, the following expression is derived: S (x) = 1.155 − 0.874x + 0.376x 2 eV, (x < 0.25) EX Si/Si0.72Ge0.18 MQW LZ = 3 nm TPL = 8.6 K NP 8000 TOSi – Si 4.3 meV 4000 TOSi – Ge TOGe – Ge TA TO + TOSi – Si 0 900 950 Photoluminescence intensity (arb units) 800 1000 900 1000 E C − E V ≈ 1.17 − 0.896x + 0.39x 2 eV, (x < 0.25) (22.13) Photon energy (meV) 1100 Si Si1–xGex Si B TO Si XNP SiGe 500 Å 500 Å 400 µm a) x = 13 % XNP SiGe XTO SiGe b) x = 17 % XNP SiGe XTO SiGe c) x = 24 % 1.6 1.5 1.4 1.3 1.2 1.1 Photon wavelength (µm) Fig. 22.21 4.2-K photoluminescence spectra from layers with the nominal structure shown in the inset (514-nm Arion laser excitation) (after Robbins et al. [22.26], copyright 1992, American Physical Society) Part C 22.2 A strain-corrected expression is also provided but this is found to modify (22.12) only slightly. Thus by adding (22.11) and (22.12) an expression for the band gap can be obtained. 1050 Energy (meV) XTO SiGe E BC (x) ≈ 0.0145 − 0.022x + 0.020x 2 eV, (x < 0.25) (22.12) 491 Counts (arb. units) (22.11) Here the presence of strain is responsible for the differences from the expression obtained for bulk samples (22.9). An expression for the exciton binding energy E BC (x) is theoretically derived for the cubic alloy and the following quadratic expression is fitted: 22.2 Optical Properties of SiGe 492 Part C Materials for Electronics 22.3 Growth of Silicon–Germanium Over the past ten years and more there have been rapid developments in techniques for the growth of Si and Si1−x Gex epitaxial layers at low temperatures. This has been made possible by a number of changes in the design of epitaxy equipment and by improvements to growth processes. There are two main prerequisites for the growth of epitaxial layers at low temperature: at 600 ◦ C at a rate of a few monolayers per second, so the hydrogen passivation approach allows epitaxial layers to be grown at low temperatures without the need for a high-temperature bake. The hydrogen-passivated surface is stable for typically 30 min after completion of the ex-situ cleaning [22.47]. • 22.3.3 Ultra-Clean Epitaxy Systems • Establishment of a clean surface prior to growth [22.44–47] Growth in an ultra-clean environment [22.48–50] The removal of oxygen and carbon is the main problem in establishing a clean surface prior to growth. A clean silicon surface is highly reactive and oxidises in air even at room temperature. The secret of low-temperature epitaxial growth is therefore the removal of this native oxide layer and the maintenance of a clean surface until epitaxy can begin. Two alternative approaches to pre-epitaxy surface cleaning have been developed, as described below. 22.3.1 In-Situ Hydrogen Bake Part C 22.3 The concept that underlies this surface clean is the controlled growth of a thin surface oxide layer, followed by its removal in the epitaxy reactor using a hydrogen bake. The controlled growth of the surface oxide layer is generally achieved using a Radio Corporation of America (RCA) clean [22.44] or a variant [22.45]. The oxide created by the RCA clean is removed in the reactor using an in-situ bake in hydrogen for around 15 min at a temperature in the range 900–950 ◦ C. The temperature required to remove the native oxide depends on the thickness of the oxide, which is determined by the severity of the surface clean. 22.3.2 Hydrogen Passivation An alternative approach to pre-epitaxy cleaning is to create an oxide-free surface using an ex-situ clean and then move quickly to epitaxial growth before the native oxide can grow. The aim of the ex-situ clean is to produce a surface that is passivated by hydrogen atoms bonded to dangling bonds from silicon atoms on the surface. When the wafers are transferred in the epitaxy reactor, the hydrogen can be released from the surface of the silicon very quickly using a low-temperature bake or even in the early stages of epitaxy without any bake. Meyerson [22.46] has reported that hydrogen desorbs Having produced a clean hydrogen-passivated silicon surface, it is clearly important to maintain the state of this surface in the epitaxy system. This necessitates the use of low-pressure epitaxy systems if epitaxial growth at low temperatures is required. Figure 22.22 summarises the partial pressures of oxygen and water vapour that need to be achieved in an epitaxy system if an oxide-free surface is to be maintained at a given temperature [22.48, 49]. This figure shows that epitaxial growth at low temperature requires low partial pressures of oxygen and water vapour, which of course can be achieved by reducing the pressure in the epitaxy system. Research [22.50] has shown that a pressure below 30 Torr is needed to achieve silicon epitaxial growth below 900 ◦ C. 22.3.4 Si1−x Gex Epitaxy The growth of Si1−x Gex epitaxial layers can be achieved over a wide range of temperatures using low-pressure chemical vapour deposition (LPCVD) [22.50] or ultra-high-vacuum chemical vapour deposition (UHVCVD) [22.51, 52]. The gas used to introduce the germanium into the layers is germane, GeH4 . The influence of germanium on the growth rate is complex, as illustrated in Fig. 22.23. At temperatures in the range 577–650 ◦ C a peak in the growth rate is seen. At low germanium contents, the growth rate increases with germanium content, whereas at high germanium content, the growth rate decreases with germanium content. In the low-temperature regime it has been proposed that hydrogen desorption from the surface is the rate-limiting step. In Si1−x Gex this occurs more easily at germanium sites than at silicon sites and hence the growth rate increases with germanium content [22.37]. As the germanium content increases, the surface contains more and more germanium and less and less hydrogen. The ratelimiting step then becomes the adsorption of germane or silane. Robbins [22.53] proposed that the sticking coefficient for germane or silane was lower at germanium sites. This would slow the adsorption rate as the ger- Silicon–Germanium: Properties, Growth and Applications Partial pressure (Torr) 1e + 0 1e – 1 H2O 1e – 2 1e – 3 O2 1e – 4 1e – 5 1e – 6 7.0 7.5 8.0 8.5 9.0 104/T (K–1) Fig. 22.22 Conditions for oxide formation in an epi- taxy system. Note that 1 atm = 1.113 bar = 760 Torr = 1.113 × 105 Pa (after Smith and Ghidini [22.48, 49], copyright Electrochemical Society 1984) Growth rate (nm/min) 25 750 °C 20 15 650 °C 10 600 °C 5 0 5 done by adding chlorine or HCl as a separate gas or by using a growth gas that contains chlorine, for example dichlorosilane, SiH2 Cl2 . With chlorine chemistry, selective growth of silicon and Si1−x Gex can be achieved to both silicon dioxide and silicon nitride. Chlorine is reported to have two effects that lead to selective growth. First it increases the surface mobility of silicon and germanium atoms, so that atoms deposited on the oxide or nitride layer are able to diffuse across the surface to the window where the growth is occurring. Second it acts as an etch [22.50] and hence can remove silicon or germanium atoms deposited on the oxide or nitride. The strength of the etching action increases with chlorine content and, if the chlorine content is too high, etching of the substrate will occur instead of epitaxial growth. A typical growth process for selective silicon epitaxy would use silane and a few percent of HCl [22.50]. The growth rate for this process is shown in Fig. 22.24, and compared with the growth rate for dichlorosilane and silane epitaxy. It can be seen that the activation energy for the silane-plus-HCl process is very similar to that for the dichlorosilane process, indicating that the growth mechanisms are similar. One disadvantage of chlorine-based growth processes over the silane process is a lower growth rate at low temperatures, as can clearly be seen in Fig. 22.24. It is also possible to grow silicon selectively using dichlorosilane and HCl [22.55]. Selective Si1−x Gex growth is generally easier to achieve than selective silicon growth, as illustrated in Growth rate (nm/min) 10 20 15 Germanium percentage (%) Fig. 22.23 Growth rate of Si1−x Gex as a function of germanium percentage for temperatures in the range 577–750 ◦ C (after Racanelli et al. [22.54], copyright 1990, American Institute of Physics) 1000 SiH4 100 DCS 10 manium content increased and hence slow the growth rate. 1 22.3.5 Selective Si1−x Gex Epitaxy Selective epitaxy is the growth of a single-crystal layer in a window, with complete suppression of growth elsewhere, and can be achieved in a number of different ways. The most common method of achieving both selective Si and Si1−x Gex epitaxy is by introducing chlorine or HCl into the growth chamber. This can either be 493 SiH4 + HCl 7 8 9 10 20 104/T (K–1) Fig. 22.24 Silicon growth rate as a function of reciprocal temperature for three different growth gases: 40 sccm of SiH4 , 80 sccm of dichlorosilane and 20 sccm of SiH4 with 2 sccm of HCl. The hydrogen flow was 2 slm (after Regolini et al. [22.50], copyright 1989 American Institute of Physics) Part C 22.3 577 °C 0 22.3 Growth of Silicon–Germanium 494 Part C Materials for Electronics Germanium percentage (%) Growth rate (nm/min) 60 1000 500 °C 550 °C 50 100 40 HCl: 10 ml / min 30 10 650 °C 600°C 20 1 10 0 Non Selective selective 0 0.2 0.4 0.6 0.8 1.0 1.2 Germane: DCS flow ratio Fig. 22.25 Germanium percentage as a function of germane: dichlorosilane (DCS) flow ratio for temperatures in the range 500–650 ◦ C showing the move from nonselective to selective growth as the proportion of germane in the gas flow increases (after Zhong et al. [22.56], copyright 1990, American Institute of Physics) Part C 22.4 Fig. 22.25 [22.56] for Si1−x Gex growth using germane and dichlorosilane. The growth moves from nonselective to selective as the proportion of germane in the gas flow increases. Arrhenius plots for Si1−x Gex growth using germane and dichlorosilane are shown in Fig. 22.26 for Si1−x Gex layers grown using germane and dichlorosilane and for two different HCl flows. It can be seen that the growth rate decreases and the activation energy increases with increasing HCl flow. The explanation proposed for this behaviour is that the limiting growth mechanism changes from hydrogen desorption from the growing surface to chlorine or HCl desorption from the surface [22.57]. This decrease in growth rate at high HCl flows is a disadvantage because it leads HCl: 40 ml / min 9.6 9.8 10.0 10.2 10.4 10.6 10.8 104/T (K–1) Fig. 22.26 Arrhenius plots for Si1−x Gex growth at two different HCl flow rates. The dichlorosilane and germane flow rates were fixed at 100 and 8 ml/min respectively (after Kiyota et al. [22.57], copyright 2002, IEEE) to increased growth times. High HCl flows can also cause surface roughening when the Si1−x Gex layer is heavily boron-doped [22.57]. These considerations demonstrate that the HCl flow should be chosen to be to the smallest value that is consistent with good selective epitaxy. Silane can be used for selective silicon epitaxy if the growth is performed at a high temperature. This approach relies on the fact that nucleation of growth on oxide is more difficult than that on silicon. This incubation time for growth on an oxide layer is relatively long at high temperatures but much shorter for growth at low temperatures. Selective silicon layers 1 µm thick can be grown using silane at a temperature of 960 ◦ C [22.58], but the achievable layer thickness decreases with decreasing temperature. At 800 ◦ C the maximum selective silicon layer thickness is around 130 nm, at 700 ◦ C it is around 60 nm, and at 620 ◦ C it is around 40 nm. Selective growth to silicon dioxide can be achieved using silane only, but not to silicon nitride. 22.4 Polycrystalline Silicon–Germanium In the past ten years there has been increasing interest in polycrystalline silicon–germanium for a number of applications that require polycrystalline material deposition at low temperature (around 600 ◦ C). Examples of potential applications are thin film transistors, gates of MOS transistors and polySiGe emitters for SiGe HBTs. In thin-film transistor technologies [22.59–61], polycrystalline silicon–germanium is compatible with the low-thermal-budget processing that is needed to produce thin-film devices for large-area electronics. The key physical property of polycrystalline silicon–germanium that makes it attractive is its lower melting point than silicon. This means that processes such as deposition, Silicon–Germanium: Properties, Growth and Applications Silicongermanium Silicon Germanium Variable 4.05 eV 4.00 eV 0.66 eV 1.12 eV Vacuum level EC EV Fig. 22.27 Band-energy levels in silicon, silicon–germanium and germanium ΦMS(Volts) 1.00 P+ poly-Si1–xGex 0.75 0.50 0.25 0.00 –0.25 N+ poly-Si1–xGex –0.50 Qf = 4 × 1010cm–2 Naub = 8 × 1014cm–3 –0.75 –1.00 0.0 0.1 0.2 0.3 0.4 Fig. 22.28 Work-function difference between a polySi1−x Gex gate and an n-type substrate as a function of germanium content (after King et al. [22.60], copyright 1994, IEEE) crystallisation, grain growth and dopant activation will occur at a lower temperature than in silicon. Thus lower temperature processes can be used for polySiGe devices and hence it is preferable to polySi in applications with tight thermal-budget requirements. In MOS transistors, polycrystalline silicon– germanium is attractive as a gate material for future generations of MOS transistor, since the germanium content in the silicon–germanium layer can be varied by 200–300 mV in the direction of a mid-gap gate [22.63–65]. This can be understood from Fig. 22.27, which compares the conduction- and valence-band energy levels in single-crystal silicon, silicon–germanium and germanium. Silicon and germanium have similar 495 electron affinities (4.05 and 4.00 eV respectively), but germanium has a much smaller band gap (0.66 eV compared with 1.12 eV). The energy difference between the valence band and the vacuum level is therefore about 0.5 eV smaller in germanium than in silicon. In silicon– germanium, this energy difference can be varied by varying the germanium content. This allows the threshold voltage of p-channel MOS transistors to be tuned by varying the germanium content in the polySi1−x Gex gate. Figure 22.28 shows values of work-function difference between the polySi1−x Gex gate and the n-type silicon substrate as a function of germanium content in a polySi1−x Gex gate [22.60]. The work function is defined as the difference in energy between the vacuum level and the Fermi level. In p+ polySi1−x Gex the Fermi level is near the valence band and hence the work-function difference varies strongly with germanium content. In n+ polySi1−x Gex the Fermi level is near the conduction band and hence the work-function difference varies little with germanium content. In Si1−x Gex HBTs, polySi1−x Gex has potential as an emitter of a SiGe HBT [22.62]. In bipolar transistors, the breakdown voltage, BVCEO , is inversely proportional to the gain [22.66] and hence transistors with a high gain have lower values of breakdown voltage. Si1−x Gex HBTs inherently have high values of the gain because the reduced band gap of the Si1−x Gex base enhances the collector current. The base current is unchanged by the Si1−x Gex base and hence the gain, which is the I(A) 10–5 Part C 22.4 0.5 0.6 Ge mole fraction 22.4 Polycrystalline Silicon–Germanium 10–6 10–7 Base, 0% Ge Base, 10% Ge Base, 19% Ge Collector 10–8 10–9 0.50 0.55 0.60 0.65 0.70 VBE(V) Fig. 22.29 Use of a polySi1−x Gex emitter to vary the base current of a Si1−x Gex HBT and hence give the best trade-off between gain and breakdown voltage BVCEO . (after Kunz et al. [22.62], copyright 2003, IEEE) 496 Part C Materials for Electronics ratio of collector current to base current, is increased. The use of a polySi1−x Gex emitter instead of a polySi emitter provides a reduced band gap in the emitter, which enhances the base current, and thereby reduces the gain. Typical measured values of base current in a polySi1−x Gex emitter are shown in Fig. 22.29, where it can be seen that 19% germanium gives a factor of approximately four reduction in gain. A polySi1−x Gex emitter therefore allows the gain to be tuned to give the best trade-off between gain and breakdown voltage BVCEO . a) 400 Minimum sheet resistance (Ω × 22.4.1 Electrical Properties of Polycrystalline Si1−x Gex Figure 22.30 shows the sheet resistance as a function of anneal temperature for boron- and phosphorus-doped Si1−x Gex for different germanium contents. For borondoped polySi1−x Gex the sheet resistance decreases with increasing germanium content, with the decrease being large between 0 and 25% germanium and smaller between 25 and 50% germanium. In contrast, for phosphorus-doped polySi1−x Gex the sheet resistance b) –1) 600 Minimum sheet resistance (Ω × 7.0 Boron doped 11.0 0.25 0.25 Si0.50Ge0.50 Si 300 11.0 2.0 100 0.25 0.25 400 200 0 500 Phosorus doped 0.25 500 300 –1) 120 7.0 0.25 200 Si0.75Ge0.25 Si0.50Ge0.50 4.0 0.25 0.25 550 600 650 700 750 Anneal temperature (°C) Si0.75Ge0.25 Si 120 7.0 0.25 550 600 650 700 750 Anneal temperature (°C) 100 0 500 Fig. 22.30 Sheet resistance as a function of anneal temperature for boron- and phosphorus-doped polycrystalline Si1−x Gex with various germanium contents (after Bang et al. [22.67], copyright 1995, American Institute of Physics) Part C 22.4 a) Percent dopant activation Hall hole mobility (cm2/Vs) 100 100 Film thickness: ~ 0.3 um Boron implant: 4 × 1015cm–2 @ 20 keV 80 Anneal: 40 min @ 900 °C 80 b) Percent dopant activation Hall electron mobility (cm2/Vs) 100 60 Film thickness: ~ 0.3 um Phosphorus implant: 4 × 1015cm–2 @ 60 keV 50 Anneal: 40 min @ 900 °C 80 40 60 60 60 30 40 40 40 20 20 0 0.0 20 0.1 0.2 0.3 0.4 0 0.5 0.6 Ge mole fraction 20 10 0 0.0 0.1 0.2 0.3 0.4 0 0.5 0.6 Ge mole fraction Fig. 22.31 Percentage dopant activation and Hall hole mobility as a function of germanium content for boron- and phosphorus-doped polySi1−x Gex (after King et al. [22.60], copyright 1994, IEEE) Silicon–Germanium: Properties, Growth and Applications increases with increasing germanium content, with the increase being small between 0 and 25% germanium and large between 25 and 50%. Similar behaviour is seen for arsenic-doped polySi1−x Gex where higher values of sheet resistance have been reported for polySi1−x Gex than for polySi [22.64]. The explanation for the sheet-resistance results in Fig. 22.30 can be found in Fig. 22.31, which shows the results of Hall measurements [22.60]. For borondoped polySi1−x Gex both the activation and the Hall mobility increase with increasing germanium content, thereby explaining the decrease in sheet resistance with References 497 increasing germanium content. For phosphorus-doped polySi1−x Gex there is little change in activation and electron mobility at low germanium contents, but a sharp decrease in activation at germanium contents above 35%. This explains the sharp increase in sheet resistance seen in Fig. 22.30 for germanium contents between 25 and 50%. The decrease in activation at high germanium contents in the phosphorus-doped polySi1−x Gex may be due to increased segregation at grain boundaries. 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